Effective Strategy for Enhancing the Performance of

Jul 4, 2019 - Anodes in Lithium-Ion Batteries: Magnetron Sputtering. Molybdenum ... ABSTRACT: The interface between the current collector and active ...
1 downloads 0 Views 8MB Size
Download PDF
Research Article Cite This: ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

www.acsami.org

Effective Strategy for Enhancing the Performance of Li4Ti5O12 Anodes in Lithium-Ion Batteries: Magnetron Sputtering Molybdenum Disulfide-Optimized Interface Architecture Xiaobo Zhu,†,‡ Xin Jiang,*,†,‡ Xiayin Yao,§ Yongxiang Leng,† Liping Wang,‡ and Qunji Xue‡

Downloaded via UNIV OF SOUTHERN INDIANA on July 19, 2019 at 07:11:17 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.

†

Key Laboratory for Advanced Technology of Materials, Ministry of Education, School of Materials Science and Engineering, Southwest Jiaotong University, Chengdu 610031, China ‡ Key Laboratory of Marine Materials and Related Technology, Zhejiang Key Laboratory of Marine Materials and Protective Technologies, Ningbo Institute of Materials Technology & Engineering, and §Ningbo Institute of Materials Technology & Engineering, Chinese Academy of Sciences, Ningbo 315201, China S Supporting Information *

ABSTRACT: The interface between the current collector and active material is the primary interface of charge transfer. Herein, we designed an effective strategy to optimize the interface architecture by depositing molybdenum disulfide on the copper foil surface (Cu−MoS2) via magnetron sputtering. The Cu−MoS2 is directly used as a current collector and supports the Li4Ti5O12 anode (Cu−MoS2−LTO). Typically, after being cycled at 1 A g−1 for 300 cycles, the capacities of the Cu−LTO cell and Cu−MoS2 cell are about 114.94 and 128.35 mA h g−1, respectively, whereas the capacity of the Cu−MoS2−LTO cell is as high as 373.9 mA h g−1 with a capacity retention rate of 89.1%. The MoS2 not only optimizes the interfacial architecture but also provides an additional capacity contribution to the Cu−MoS2−LTO cell. Based on scanning electron microscopy and X-ray photoelectron spectroscopy test analysis, we propose a dual interface model. It is revealed that the molybdenum disulfide film can significantly improve the charge-transfer efficiency and uniformity of the interface, reduce internal resistance of the batteries, prevent oxidation of the copper foil, and thereby improve the chemical stability of the current collector. In addition, magnetron sputtering technology has large-scale productivity and greatly enhances the industrial application of this strategy. KEYWORDS: molybdenum disulfide, magnetron sputtering, interface modification, Cu−MoS2 current collector, dual interface model

1. INTRODUCTION Lithium-ion batteries (LIBs) as a clean energy source are one of the main power sources for future hybrid electric vehicles and community energy because of their high energy density and high operating voltage.1−5 Nowadays, with the gradual deepening of the electrode interface research, researchers have gradually paid attention to the influence of the current collector on the performance of the electrode, such as rate performance and cycle life.6,7 The current collector not only functions to carry the active material but also collects the electrons produced by the electrochemical reaction and conducts it to the external circuit, thereby realizing the process of converting chemical energy into electrical energy. In the conventional LIB electrode manufacturing process, the electrode slurry is directly uniformly coated on the surface of the current collector, and the bonding between the active material and current collector is achieved by the binder. However, the electrode manufacturing process has the following two drawbacks that cause battery cycle performance and safety to decrease: (1) the contact area between the rigid collector surface and the active material is limited, resulting in an increase in the resistance of interface, especially at high © XXXX American Chemical Society

current densities. (2) The bonding strength of the binder is limited, and the active material is easily detached from the surface of the current collector during charge−discharge, which causes the internal resistance of the LIBs to further increase. At present, the methods for solving the above problems mainly include increasing the amount of binder, corroding, and roughening of the current collector surface, coating a conductive coating, and so forth, and the purpose thereof is to improve the bonding strength between the active material and current collector.8−11 However, the binder itself is an insulator. Increasing the amount of binder not only increases the overall internal resistance but also may cause the heat generation of the lithium-ion capacitor to rise, posing a safety hazard. Moreover, the bond between the conductive additives and current collector is mainly a physical action, which inevitably leads to a poor bonding force between the conductive additives and current collector. Although the corrosion treatment can coarsen the surface of the current Received: April 25, 2019 Accepted: July 4, 2019 Published: July 4, 2019 A

DOI: 10.1021/acsami.9b07269 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces

Figure 1. Cu−MoS2 current collector preparation process and digital photograph (135 × 300 mm2).

of the interface between the active material and the current collector was not revealed. Moreover, the above method also has problems in that it is difficult to realize the large-scale safe production of full battery components. Fortunately, the magnetron sputtering technology can effectively solve the above problems because of its mature technology, low deposition temperature, large-area preparation, and strong film-based bonding.21 At present, our laboratory can produce Cu−MoS2 current collectors with a size of 40 cm × 60 cm at a time. Moreover, continuous coating on flexible substrates using roll-to-roll magnetron sputtering technology has achieved industrialization. In this work, we have designed an effective strategy to optimize the interface architecture between the current collector and active materials. This strategy is achieved by depositing a layer of molybdenum disulfide conductive layer on the copper foil surface (Cu−MoS2) via magnetron sputtering. To demonstrate the effectiveness of this strategy, we discussed the difference in interface properties between the two current collectors and studied the effect of MoS2-modified layer on the cell performance. The MoS2 not only optimizes the interfacial structure between the current collector and LTO anode material, but also provides an additional capacity contribution to the Cu−MoS2−LTO cell. Based on scanning electron microscopy (SEM) and X-ray photoelectron spectroscopy (XPS) test analysis, a dual interface model was proposed to illustrate the transport mechanism of Li+ and electrons. The MoS2 film can significantly improve the chargetransfer efficiency and uniformity of the interface, prevent oxidation of the copper foil, and thereby improve the chemical stability of the current collector.

collector and improve the bond strength between the active material and the current collector, the corrosion process will cause certain surface damage to the current collector and reduce the mechanical strength of the current collector, which will seriously affect the subsequent finished product processing. Therefore, the most effective method at present is to apply a conductive coating on the surface of the current collector. Molybdenum disulfide (MoS2), a typical two-dimensional nanomaterial, has attracted extensive attention in LIBs because of its weak van der Waals force, large inter-layer spacing (0.615 nm), and high theoretical lithium storage capacity.12,13 Herein, this study used MoS2 as a conductive coating to enhance the interfacial properties of the current collector. The method of coating the conductive coating on the surface of the current collector mainly includes electrochemical etching,10,14 chemical vapor deposition (CVD),7,15 and chemically oxidized processes.16,17 Kim and Choi18 prepared a modified Cu current collector via depositing a thin layer of graphene film on the copper foil surface by CVD. However, the CVD technology requires a high deposition temperature, consumes a lot of energy, and is not suitable for a low melting point current collector such as aluminum. Wang et al.19 used an oxide coated on the Al current collector surface to enhance the corrosion resistance of the current collector, but this modification process requires many chemical reagents such as NaOH, Na2CO3, KMnO4, HNO3, sodium dodecyl sulfate, potassium fluozirconate, and so forth, and the preparation process is complicated. Simon et al.20 modified the nanoarchitectured Cu foil surface by electrochemically assisted template and electrodeposition techniques and obtained a battery with good performance, but the expensive anodic aluminum oxide template limits its mass production. Wang et al.6 made a discharge method that improved the surface properties of the Cu foil by a carbon layer, but the mechanism B

DOI: 10.1021/acsami.9b07269 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces

Figure 2. (a) Raman spectra, (b) XPS Mo 3d spectrum, and (c) XPS S 2p spectrum of the MoS2 film. angle of pure Cu and Cu−MoS2 films were measured at ambient temperature using a measuring apparatus (OCA20).

2. EXPERIMENTAL SECTION 2.1. Batch Preparation of Cu−MoS2 Current Collector Film. The preparation process of the Cu−MoS2 current collector is shown in Figure 1. A layer of the MoS2 film was deposited on the copper foil (wt % ≥99.8%) with a size of 9 μm × 13.5 cm × 30 cm by magnetron sputtering. Before deposition, the Cu foil was swabbed successively with acetone and ethanol and then loaded into a vacuum chamber of the multifunctional magnetron sputtering equipment (Teer PlasMag CF-800). The coating process is started when the pressure in the vacuum chamber is lower than 5.0 × 10−4 Pa. Prior to deposition of the MoS2 film, Ar+ bombardment was used for 20 min under Ar of 35 sccm and −50 V bias voltage to remove impurities on the Cu foil. During the deposition, the Ar flow rate and bias were kept constant, and the MoS2 film with a thickness of about 2.3 μm was deposited by sputtering a MoS2 target (purity 99.99 at. %) for 4 h. 2.2. Batch Fabrication of Electrodes. The pure Cu foil and Cu−MoS2 were directly used as the current collector and support spinel Li4Ti5O12 (without any treatment, Supporting Information, Figure S1) anode. Super P (carbon black) and polyvinylidene fluoride (PVDF) were used as a conductive agent and binder, respectively. The slurry was prepared by mixing LTO/Super P/PVDF in a weight ratio of 8:1:1. The anode slurry was then stirred by centrifugation at room temperature for 24 h. The mixed slurry was separately applied to the two current collectors by blade coating machine and then dried at 105 °C for 12 h and punched into wafers with a diameter of 13 mm. In addition, to determine the capacity contribution of MoS2, the Cu−MoS2 film directly punched into wafers with a diameter of 13 mm as the anode electrodes. Among them, the effective mass of anode materials of the Cu−LTO cell and Cu−MoS2 cell are 3 and 1.45 mg, respectively, and the effective mass of the anode material of Cu− MoS2−LTO cell is the superposition of two anode electrode materials (4.45 mg). 2.3. Fabrication and Performance Test of Batteries. The Cu−LTO, Cu−MoS2, and Cu−MoS2−LTO were used as working electrodes, polypropylene (thickness is 25 μm) film was used as a separator and lithium metal (Φ 16 mm × 0.6 mm, wt % ≥99.9%) was used as a counter electrode to assemble CR-2032 button cells. The electrolytes were a 1 M LiPF6 consisting of 1:1:1 ethylene carbonate, dimethyl carbonate, and diethyl carbonate. The cells were performed to a galvanostatic charge−discharge experiments at different current densities using a CT2001A battery testing system (LAND Electronic Co.) at a voltage of 0.01−3.0 V (vs Li/Li+). The electrochemical workstation (CHI660E) was used to measure the electrochemical impedance spectroscopy (EIS) and charge−discharge before and after the cell cycle. The test frequency range was set to 10−2 to 105 Hz, the amplitude was controlled at 5 mV, and the test voltage interval is 0.01−3.0 V. After charge−discharge cycling, the electrodes in the cell were extracted in a glove box, then washed with acetone, and vacuumdried for morphological characterization. 2.4. Characterizations. The structure of MoS2 films was measured using a Renishaw inVia spectrometer (Renishaw, England) with a laser selection of 514.5 nm Ar+. An X-ray photoelectron spectrometer (Axis Ultra DLD, Kratos, England) was used for measuring the chemical composition of materials. The microstructure of the materials were characterized by S4800 (Hitachi) and a FEI Quanta FEG 250 (FEI) scanning electron microscope. The contact

3. RESULTS AND DISCUSSION Figure 2 gives the Raman and XPS spectra of the MoS2 film, respectively, which can effectively characterize the composition and structure of the conductive coating. There are two characteristic peaks in the Raman spectra (Figure 2a), including the Mo−S in-plane E12g vibration mode at 378.6 cm−1 and Mo−S out-of-plane A1g vibration mode around 405.6 cm−1,22,23 indicating that the conductive coating deposited by magnetron sputtering is a typical amorphous MoS2 film. Moreover, Figure 2b,c shows the Mo 3d and S 2p spectrum of MoS2 film under Gaussian fitting. For the Mo 3d, the peaks around 225.9, 229.6, and 231.8 eV correspond to S 2s, Mo 3d5/2, and Mo 3d3/2 of MoS2, respectively.24 For the S 2p, the binding energies of S 2p3/2 and S 2p1/2 are located at 162.3 and 163.4 eV, respectively.25 It is shown that the film deposited by magnetron sputtering is a pure MoS2 film. Figure 3a,b shows the surface morphology and roughness of pure Cu foil and MoS2 film, respectively. Obviously, the MoS2

Figure 3. Surface morphology and contact angle: (a,c) pure copper foil and (b,d) Cu−MoS2 film.

film deposited on the surface of the Cu foil is very uniform, exhibiting a micro/nanostructure with a rough surface. This rough surface can increase the contact area between the Cu− MoS2 current collector and active material, thereby increasing the bond strength and electron transport rate between them, according to the semiconductor-electrochemical model (Figure S2).26 When the metal current collector or the conductive coating with good conductivity support the active material, a typical Schottky or Schottky-like structure is formed, and the chemical potential of the electrolyte is equivalent to the Fermi level of the electrolyte. When the initial system is stabilized, the Fermi level is in equilibrium, and there is no transfer or C

DOI: 10.1021/acsami.9b07269 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces reaction of electrons and ions. When a negative voltage is applied, Esub is shifted to the conduction band level, and the electrochemical reaction proceeds at the interface between the active material and electrolyte. The electron transfer occurs at the contact surface of the active material and MoS2. According to the point charge potential equation (eq 1), the charge of the contact surface should be equal to the product of the current density and contact area (eq 2). Therefore, the relationship equation of the contact potential ηsub can be obtained (eq 3). η = kQ /r

(1)

η is the potential of the point charge, k is the electrostatic constant, Q is the amount of charge, and r is the distance between the charge and the potential. Q s = isSs

(2)

ηsub = kisSs/r

(3)

Figure 4. (a) EIS curves and (b) equivalent circuit model diagram of Cu−LTO, Cu−MoS2, and Cu−MoS2−LTO cells after 1 cycle at 1 A g−1.

LTO cell after one charge−discharge cycle. The Nyquist plots (Figure 4a) of the three cells both contain a high-frequency region (the ionic resistance of the electrolyte, Rs), semicircular mid-frequency region (resistance of the interface between the current collector and the LTO, Ril), and line at the lowfrequency range (Warburg impedance, Zw).29 Obviously, the Cu−MoS2−LTO cell has the lowest total impedance, indicating that the total resistance of the Cu−MoS2−LTO cell is the smallest. The equivalent circuit model diagram is shown in Figure 4b, where the CPE represents the total capacitance of the passivation film and the electric double layer. After one charge−discharge cycle, the Rs of Cu−LTO cell (3.89 Ω), Cu−MoS2 cell (3.98 Ω), and Cu−MoS2−LTO cell (3.77 Ω) are not much different, which is because the electrolyte composition in the initial stage is the same. However, the Ril of Cu−LTO cells, Cu−MoS2 cells, and Cu−MoS2−LTO cells are 21.81, 39.44, and 18.56 Ω, respectively. The Ril of the Cu−MoS2−LTO cell is more similar to that of the Cu−LTO cell, indicating that the Ril of the Cu−MoS2−LTO cell represents the resistance of the interface between the Cu−MoS2 current collector and LTO, that is, the current collector in the Cu−MoS2−LTO cell is Cu−MoS2. In addition, the Ril of the Cu−MoS2−LTO cell is smaller than that of the Cu−LTO cell. This can be ascribed to the fact that the rough MoS2 film can promote the charge transfer and reduce the resistance of the interface between the Cu−MoS2 and the LTO, thus enhancing electron collection capability and improving electrochemical kinetics. To calculate the capacity contribution of the MoS2 film to the Cu−MoS2−LTO cell, we tested the charge−discharge performance of the Cu−MoS2 cell, Cu−LTO cell, and Cu− MoS2−LTO cell, respectively. Figure 5a displays the typical Cu−LTO cell galvanostatic charge−discharge curves of the initial three cycles under 1 A g−1. A typical charge−discharge platform appears near 1.55 V, and its polarization voltage (ΔV, the voltage difference between the discharge platform and charge platform) is about 0.059 V. Figure 5b shows the rate capabilities of the Cu−LTO cell at different current rates. At a current density of 0.1, 0.2, 0.5, 1, and 2 A g−1, the Cu−LTO cell discharge specific capacity of 159.4, 149.1, 139.3, 127.9, and 115.07 mA h g−1, respectively, and the reversible capacity retention was approximately 94.79%. The Cu−LTO cell has poor charge and discharge performance, especially at high current densities (2 A g−1). As can also be seen from Figure 5c, the Cu−LTO cell has low capacity and ordinary cycle stability. The initial discharge specific capacity and Coulombic efficiency (CE) of Cu−LTO cell are about 128.96 mA h g−1 and 92.87%, respectively. Over the next 100 cycles, the discharge specific

Qs is the amount of charge between the active material and MoS2, is represents the current density, and Ss is the contact area. Obviously, the degree of the Ef offset is directly controlled by ηsub; increasing the contact area Ss can increase the driving force ηsub of the chemical reaction, causing Ef to shift to the depth of the conduction band level. Therefore, the larger the contact area of the current collector with the active material, the smaller the degree of polarization of the electrode is, resulting in a significant increase in the rate performance of the battery. To better illustrate the difference in surface characteristics between the pure Cu foil and the Cu−MoS2 current collector, we use the electrolyte as the test medium to test the contact angles of the two current collectors (Figure 3c,d). The contact angle of pure Cu foil is approximately 85.2°, whereas the contact angle of the Cu−MoS2 current collector surface is increased to 135.7° because of the presence of the micro/ nanostructure, exhibiting poor wettability. In general, the poorer the wettability of the film, the better the corrosion resistance is, because the film can effectively reduce the contact between the corrosive medium and substrate.27 Thus, the corrosion medium cannot penetrate into the film.28 In other words, the MoS2 film can effectively slow down the corrosion of the electrolyte to the current collector, increase the stability of the Cu−MoS2, and extend the service life of the current collector. In addition, the bonding strength of the MoS2 film deposited via magnetron sputtering to the copper foil is good (Figure S3). The adhesion cross-cut test shows that the bonding force between LTO and Cu−MoS2 is greater than that between LTO and pure Cu, indicating that the MoS2 film can enhance the adhesion between the Cu−MoS2 and LTO, which is beneficial to improve the performance of the cell and extend the cell service life. The strong bonding between the MoS2 film and the pure copper foil lays a good foundation for the reliable application of the robust Cu−MoS2 current collector. To demonstrate the effectiveness of this strategy for the modification of the interfacial properties between the current collector and the active material, we performed a series of battery performance tests including EIS test, galvanostatic charge−discharge test, discharge rate capabilities test, cycling performances test, and cyclic voltammetry (CV) test for the Cu−LTO cell, Cu−MoS2 cell, and Cu−MoS2−LTO cell. Figure 4 displays the EIS curves and equivalent circuit model diagram of the Cu−LTO cell, Cu−MoS2 cell, and Cu−MoS2− D

DOI: 10.1021/acsami.9b07269 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces

Figure 5. Performance of the Cu−LTO cell: (a) charge−discharge profiles of the initial three cycles under 1 A g−1, (b) rate capability, and (c) cycling performance under 1 A g−1.

Figure 6. Performance of Cu−MoS2 cell: (a) charge−discharge profiles of the initial three cycles under 0.1 A g−1, (b) rate capability, and (c) cycling performance under 1 A g−1.

capacity stabilize at 124.29 mA h g−1 and the capacity retention is 96.38%. Subsequently, the performance of the Cu−LTO cell began to decline. At 300 cycles, the discharge specific capacity and the capacity retention are approximately 114.94 mA h g−1 and 89.1%, respectively. Figure 6a presents the typical Cu−MoS2 cell galvanostatic charge−discharge curves of the initial three cycles under 0.1 A g−1. The Cu−MoS2 cell presents high initial discharge and charge capacities of 737.84 and 606.39 mA h g−1, respectively. The corresponding CE is 82.18%. The capacity consumption of the Cu−MoS2 cell is mainly the formation of a solid electrolyte interphase (SEI) during the first charge−discharge. Figure 6b shows the rate capabilities of Cu−MoS2 cell at different current rates. It can be clearly seen that the Cu−MoS2

cell has poor stability, poor cycle performance, and small capacity at a large current density. The discharge specific capacity of the Cu−MoS2 cell is 634.1, 486.9, 323.4, 253.2, and 185.8 mA h g−1 at a current density of 0.1, 0.2, 0.5, 1, and 2 A g−1, respectively. In addition, the reversible capacity retention is merely 53.21%. Likewise, similar conclusions can be drawn from the cycle performance of the Cu−MoS2 cell under 1 A g−1 (Figure 6c). After 100 cycles, the discharge specific capacity rapidly drops to 280.18 mA h g−1, which is less than half of the initial discharge capacity. After 300 cycles, the discharge specific capacity was only 128.35 mA h g−1, and the capacity retention rate was approximately 23.2%. However, it is undeniable that the contribution of MoS2 to the capacity of the cell cannot be ignored. However, from the charge−discharge E

DOI: 10.1021/acsami.9b07269 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces

Figure 7. Performance of Cu−MoS2−LTO cell: (a) charge−discharge profiles of the initial three cycles under 1 A g−1, (b) rate capability, and (c) cycling performance under 1 A g−1.

cycle test of Cu−MoS2−LTO cell (Figure 7), it can be seen that the Cu−MoS2−LTO cell has a high charge−discharge ratio capacity and excellent cycle stability, which can also prove the effectiveness of this strategy for optimizing the interface performance of the current collector. Figure 7a exhibits the typical Cu−MoS2−LTO cell galvanostatic charge−discharge curves of the initial three cycles under 1 A g−1. The charge and discharge platform of the Cu−MoS2−LTO cell includes a typical charge and discharge platform of MoS2 and Li4Ti5O12, and its polarization voltage is about 0.029 V. The polarization voltage of Cu−MoS2−LTO cell is smaller than that of the Cu−LTO cell, indicating that the MoS2 film can effectively decrease the polarization of the cell. It may be due to better electronic contact between Cu− MoS2 and the active material, which enable electron transfer to perform faster. This is consistent with the EIS analysis. In addition, the Cu−MoS2−LTO cell presents high initial discharge and charge capacities of 457.07 and 441.55 mA h g−1, respectively, corresponding to a large CE of approximately 96.6%. Figure 7b presents the rate capabilities of Cu−MoS2− LTO cell at different current rates. At a current density of 0.1, 0.2, 0.5, 1, and 2 A g−1, the Cu−MoS2−LTO discharge-specific capacity was 659.2, 595.0, 489.1, 430.0, and 355.3 mA h g−1, respectively, and the reversible capacity retention was approximately 82.77%. Cu−MoS2−LTO cells still exhibit high specific capacity and good stability even when charged and discharged at high current densities. Besides, the Cu− MoS2−LTO cell exhibited longer cycle life performance at higher current rates of 1 A g−1 (Figure 7c). The initial discharge specific capacity of the Cu−MoS2−LTO cell at 1 A g−1 is approximately 422.3 mA h g−1, after 100 cycles, the capacity retention rate is as high as 95.4%, and after 300 cycles, the capacity retention rate is still as high as 89.1%. It shows that the Cu−MoS2−LTO cell has a high specific capacity and excellent cycle stability. According to the above data analysis, the deposition of the MoS2 film on the surface of copper foil by magnetron sputtering technology can effectively improve the interface performance between the current collector and active material.

At the same time, MoS2 also provides capacity contribution to Cu−MoS2−LTO cells as an additional anode material. Specifically, the effective mass of anode materials in Cu− LTO, Cu−MoS2, and Cu−MoS2−LTO cells are 3.00, 1.45, and 4.45 mg, respectively, and the specific discharge capacity after one charge−discharge cycle at 1 A g−1 is 129.28, 554.57, and 422.3 mA h g−1, respectively. If the specific capacity of Cu−MoS2−LTO cell is simply superposition calculation, the theoretical value of discharge specific capacity after one charge−discharge cycle in 1 A g−1 should be 267.86 mA h g−1. Therefore, the improvement of 154.44 mA h g−1 discharge specific capacity of the Cu−MoS2−LTO cell can be attributed to the interface modification of this strategy. After 300 cycles, the specific discharge capacity of Cu−LTO, Cu−MoS2, Cu− MoS2−LTO cells are 114.94, 128.35, and 376.3 mA h g−1, respectively. If the specific capacity of the Cu−MoS2−LTO cell is simply superposition calculation, the theoretical value of discharge specific capacity after 300 charge−discharge cycles in 1 A g−1 should be 119.31 mA h g−1. Although MoS2 has a capacity contribution, its stability is poor, particularly at 300 cycles. This indicates that the MoS2 optimization interface structure between the collector and anode material is the most important factor to improve the overall performance of the cell. This structure can effectively slow down the capacity decay of LTO and MoS2 in Cu−MoS2−LTO cells. Hence, the Cu− MoS2−LTO cell can combine the high capacity of the MoS2 with the stability of the LTO to exhibit high capacity and excellent stability. This is attributed to the fact that the rough Cu−MoS2 surface can increase the contact area between the current collector and active material and reduce the resistance of the interface, thereby reducing the resistance of the Cu− MoS2−LTO cell and improving the stability of the MoS2. To further unveil the high rate property of the Cu−MoS2− LTO cell in LIBs, the CV curves of the Cu−MoS2 cell of the first three cycles at a scan rate of 0.1 mV s−1 in a potential range of 0.01−3.0 V (vs Li/Li+) are shown in Figure S4. In addition, the CV curves of the Cu−MoS2−LTO cell at different sweep rates in the potential range of 0.01−3.0 V are collected in Figure 8a. When the sweep speed is 1.0 mV s−1, F

DOI: 10.1021/acsami.9b07269 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces

Figure 8. (a) Cyclic voltammogram profiles at different sweep rates and (b) capacitive contribution to the total capacity at various scan rates of Cu−MoS2−LTO cell.

the Cu−MoS2−LTO cell showed two reduction peaks in the cathodic process. The first peak located at 1.22 V is due to the conversion reaction of LTO (two-phase reactions between the spinel phase structured LTO and the rock salt phase structure of Li7Ti5O12, eq 4) and the phase transformation of MoS2;6 the phase transformation is due to the intercalation of Li+ into MoS2 to convert MoS2 into LixMoS2.30 The second peak located at 0.48 V is attributed to the conversion reaction of LixMoS2 (eq 5).31 Li4Ti5O12 + 3Li+ + 3e− → Li 7Ti5O12

(4) Figure 9. (a) EIS curves and (b) equivalent circuit model diagram of Cu−LTO, Cu−MoS2 and Cu−MoS2−LTO cells after cycling for 300 cycles at 1 A g−1.

LixMoS2 + (4 − x)Li+ + (4 − x)e− → Mo + 2Li 2S (5)

In the anodic process, the Cu−MoS2−LTO cell showed two oxidation peaks. First, Li7Ti5O12 is converted to LTO (eq 6), and the first oxidation peak appears near 1.84 V.32 Then, the Li+ stored on Mo nanoparticles begin to be released, the Mo nanoparticles are partially oxidized to MoS2 and Li2S is oxidized to S (eq 7), and the second oxidation peak appears in the vicinity of 2.35 V.33 Li 7Ti5O12 → Li4Ti5O12 + 3Li+ + 3e−

(6)

Li 2S ↔ S + 2Li+ + 2e−

(7)

show a smaller electrochemical impedance after a long charge− discharge cycles. This is caused by the disruption of the passive film on the surface of the lithium metal. The equivalent circuit model diagram is shown in Figure 9b, where the RSEI represents the resistance of the SEI formed during charge− discharge. After 300 cycles, the Rs of all three cells is increased, mainly because during the charge−discharge test of the cell, the electrolyte generates impurities during the redox reaction, thereby increasing the electrolyte resistance Rs. Specifically, the Rs of the Cu−MoS2−LTO cell is 8.05 Ω, which is smaller than that of the Cu−LTO cell (14.07 Ω for Rs). This can be ascribed to that the MoS2 film of the Cu−MoS2−LTO cell can increase the contact area between the current collector and negative electrode material, which can increase the electron transport and reduce the Li+ diffusion, thereby reducing the formation of impurities in the electrolyte. However, the Rs and Ril of Cu−MoS2 cell is about 15.88 and 16.47 Ω, respectively, which are associated with the destruction of the MoS2 structure during charging−discharging. It is also the reason for the sharp decline in the stability of the Cu−MoS2 cell. In addition, the RSEI of the Cu−LTO cell and Cu−MoS2−LTO cell is 7.78 and 4.14 Ω, respectively. This may be because the SEI in the Cu−MoS2−LTO cell is more uniform and thinner and has less impurity, which can be attributed to that the existence of the MoS2 film can effectively slow the formation of the SEI. Similarly, after 300 cycles, the Ril of the Cu−MoS2− LTO cell is about 5.43 Ω, which is also smaller than that of the Cu−LTO cell (8.69 Ω for Ril). This is mainly due to the rapid conduction of electrons and Li+ at the interface during charge and discharge of the cell, resulting in a decrease in the resistance of the interface. The morphologies (S4800, Hitachi) of the Cu−LTO electrode and Cu−MoS2−LTO electrode after charge− discharge 300 cycles at a current rate of 1 A g−1 are shown

It is assumed that the peak current value (i) and sweep rate (ν) in the CV measurement have the following relationship: i = aνb, in which a and b are variables. The b value is the slope of the log(i) versus log(ν) curve, and the b values of the four peaks are 0.89, 0.76, 0.92, and 0.68, respectively. All of the b values are greater than 1/2, indicating that the electrochemical process of the Cu−MoS2−LTO cell is dominated by capacitance.34,35 At a certain potential, the current (i) can be described as i = k1ν + k2ν1/2, where k1ν and k2ν1/2 represent the capacitive effect and the diffusion-dominated reaction, respectively.36 k1 and k2 are the slope and intercept of the equation, representing the contribution of the capacitance and contribution of the diffusion, respectively. As the scan rate increases, the proportion of capacitance contribution increases gradually. The capacitance contribution of the Cu−MoS2− LTO cell at 0.2−1 mV s−1 is 65.77, 69.32, 73.48, 79.56, and 86.74%, respectively (Figure 8b). This indicates that the diffusion-control reactions of the Cu−MoS2−LTO cell are restrained, which may be the reason for the cell’s ability to achieve high rates. Figure 9 reveals the EIS curves and equivalent circuit model diagram of the Cu−LTO cell, Cu−MoS2 cell, and Cu−MoS2− LTO cell after cycling for 300 cycles at 1 A g−1. All three cells G

DOI: 10.1021/acsami.9b07269 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces

electrode remained intact (Figure 10f), and the LTO was firmly bonded to the Cu−MoS2 current collector. This strong adhesion promotes good electrical contact between the Cu− MoS2 current collector and LTO, which is the main reason why the Cu−MoS2−LTO cell has superior rate performance and excellent cycle performance. However, it can be seen from the cycle performance test of the Cu−MoS2 cell that although MoS2 has a capacity contribution, its stability is poor, particularly at 300 cycles (Figure 6). Therefore, the capacity contribution of the MoS2 in the Cu−MoS2−LTO cell will decrease with the charging−discharging progresses. Therefore, although the structure of the Cu−MoS2−LTO cell is more stable than that of the Cu−LTO cell, the capacity retention of the Cu−MoS2−LTO cell is not further improved because of the capacity attenuation of the MoS2. In addition, Figure S6 displays the SEM images of the Cu−MoS2 electrode before and after the cycling. The results also show that in the Cu− MoS2−LTO cell, on the one hand, the rapid transfer of Li+ and electrons at the interface between LTO and MoS2 can reduce the thermal fatigue of the Cu−MoS2 current collector. On the other hand, the surface of the MoS2 is in close contact with the LTO. Thereby effectively buffering the volume expansion of MoS2, the Cu−MoS2−LTO cell can still maintain excellent cycle performance. To determine the distribution of the SEI in the cells, we performed an energy spectrum analysis on the cross-sectional of the Cu−LTO electrode and Cu−MoS2−LTO electrode. Figure 11 is the SEM (FEI Quanta FEG 250) cross-sectional

in Figure 10. It can be clearly seen from Figure 10a that a thin film covers the surface of the Cu−LTO electrode. This is due

Figure 10. SEM images of the Cu−LTO electrode: (a) low magnification, (c) high magnification, (e) cross-sectional, and Cu− MoS2−LTO electrode: (b) low magnification, (d) high magnification, (f) cross-sectional after 300 cycles at 1 A g−1.

to the decomposition of the electrolyte during charge− discharge and the SEI formed on the surface of the electrode, which can effectively prevent the passage of solvent molecules, but Li+ can freely do intercalation−deintercalation through the passivation layer. However, the SEI on the surface of the Cu− LTO electrode is thick and uneven and even blocks the porous structure of the LTO on the surface, which is disadvantageous for electrolyte penetration and hinders Li+ transport. This is more pronounced at high magnification (Figure 10c, Figure S4), and it can be seen that the LTO on the surface of the Cu− LTO electrode is completely covered by the SEI. Furthermore, the inset in Figure 10a is an optical photograph of the surface of the Cu−LTO electrode after the cycle. It can be clearly seen that the LTO on the pure Cu surface has peeled off partly, which is the result of reduplicative Li+ intercalation− deintercalation. The destruction of the electrode structure hinders electron transport and Li+ diffusion, inevitably causing capacity decay. In contrast, the surface of the Cu−MoS2−LTO remains intact (Figure 10b). It can be seen from the inserted image that the electrode slurry remains firmly attached to the surface of the Cu−MoS2 current collector after the cell is cycled 300 times at a current density of 1 A g−1, indicating that the adhesion of Cu−MoS2−LTO between the current collector and active material is better than that of Cu−LTO. At high magnification (Figure 10d), the Cu−MoS2−LTO electrode still maintains a complete porous structure and is covered with a very thin SEI on the surface, which helps prevent impurities in the electrolyte from penetrating into the LTO but can provide a transportation channel for Li+. Furthermore, it can be seen from the cross-sectional image of the Cu−LTO electrode (Figure 10e) that there are many distinct cracks in the active material, which are caused by polarization during charging−discharging. In contrast, the structure of the active material in the Cu−MoS2−LTO

Figure 11. SEM cross-sectional images and EDS element distribution of the electrodes after cycling: (a) Cu−LTO electrode; (b) Cu− MoS2−LTO electrode.

images and energy-dispersive spectroscopy (EDS) element distribution of two electrodes after cycling. To remove the electrolyte on the surface of the electrodes, the electrodes were sequentially ultrasonically washed in acetone and ethanol for 15 min. However, the ultrasonic process inevitably causes mechanical damage to the bonding force of the electrodes, resulting in a significant separation of the LTO from the current collector. It can be seen from the EDS element distribution diagram of Cu−MoS2 electrode (Figure 11a) that the surface of LTO, electrode crack, and interface between the H

DOI: 10.1021/acsami.9b07269 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces

pure Cu current collector, including Cu 2p3/2 (932.8 eV), Cu 2p1/2 (952.7 eV), and two weak satellite peaks Cu(II) at 942.8 and 962.7 eV, which is assigned to typical Cu 2p spectra of CuO (Figure 12a).6 It is indicated that trace amounts of copper oxide were generated on the pure Cu surface after a long cycling charge−discharge process. In addition, Taniguchi and Kalimuldina reported that Li2S and Cu will form CuxS in the vicinity of the oxidation process potential of 2.25 V.37−39 However, the XPS spectra of Cu−MoS2 surface contains only copper phases: Cu 2p3/2 at 932.8 eV and Cu 2p1/2 at 952.7 eV (Figure 12b), revealing that the surface of the copper foil is not oxidized to CuO or CuxS because of the presence of the MoS2 film. Furthermore, the MoS2 film on the Cu−MoS2 electrode after the cycle is still coated on the copper foil surface because of the strong adhesion between the MoS2 film and copper, which is exactly an inherent advantage of magnetrons sputtering technology. There are two distinct peaks in Figure 12c, corresponding to Mo 3d3/2 of MoS2 and Mo 3d3/2 of MoO3, respectively.40 Similarly, the three peaks in Figure 12d correspond to S 2p3/2 and S 2p1/2 of MoS2 and S 2p1/2 of SO2, respectively.41 Compared with the initial MoS2 spectrum, the MoS2 film was significantly oxidized after charge−discharge cycling. It is this self-oxidation of the MoS2 film that protects the copper foil from oxidation. This benefits stabilize cell performance and extend cycle life. Therefore, the robust MoS2 film can be used as a protective film and maintain Cu−MoS2 current collector integrity and stability throughout the long cycles. Based on the above discussion, a dual interface model was proposed to explain this electrochemical difference by detailed analysis of the charge-transfer path and mode of the electrode. The dual interface refers to the SEI interface between the current collector and anode material and SEI interface on the electrode surface, which has been confirmed by SEM in Figure 11. The schematic diagram of the charge transport mechanism at the dual interface is shown in Figure 13. During the charging−discharging process, a thin SEI is gradually formed on the surface of the electrode, and this SEI can effectively block the solvent molecules in the electrolyte and allows electrons and Li+ to pass quickly. However, meanwhile, the cell during charging−discharging is exothermic, resulting in the

LTO and pure Cu current collector are the enrichment area of F element, indicating that the SEI exists on the surface and crack of the Cu−MoS2 electrode and the interface between pure Cu current collector and LTO. However, as can be seen from Figure 11b, the SEI in the Cu−MoS2−LTO electrode exists at the surface of the Cu−MoS2−LTO electrode and the interface between the LTO and the MoS2 film. Most importantly, the SEI in the Cu−MoS2−LTO electrode is thinner and more uniform. It is indicates that the MoS2 film can inhibit the formation of the SEI and significantly improve the charge-transfer efficiency and uniformity of the interface. To study the chemical stability of the current collector in the electrolyte, we used XPS to detect the microscopic composition of the pure Cu and Cu−MoS2 current collector after 300 charge−discharge tests, as shown in Figure 12. The

Figure 12. XPS spectra: (a) Cu 2p of pure Cu current collector after 300 cycles, (b) Cu 2p, (c) Mo 3d, and (d) S 2p of Cu−MoS2 current collector after 300 cycles.

LTO on the pure Cu current collector and the Cu−MoS2 current collector surface was wiped off with ethanol. Obviously, there were four peaks in the Cu 2p spectra of

Figure 13. Schematic diagram of the charge transport mechanism at the dual interface: (a) Cu−LTO cell and (b) Cu−MoS2−LTO cell. I

DOI: 10.1021/acsami.9b07269 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces

applicable to commercial carbon-based anode materials or advanced next-generation Si−C anodes and can also be used to improve interfacial properties between the Al current collector and cathode materials.

thermal effect. For the Cu−LTO cell, the thermal effect not only promotes the formation of high resistance oxide layer on the Cu current collector surface, hindering electron transmission but also produces thermal strain between the current collector and the active material and reduces the bonding strength between them. On the other hand, alternating temperature changes can cause thermal fatigue in the current collector. The thermal fatigue can create cracks between the current collector and the active material to weaken the interface bonding, and at the same time, the electrolyte penetrates into the void and is concentrated in the weak binding region of the interface between the pure Cu current collector and LTO to form an SEI (Figure 13a). However, as the number of cycle’s increases, the SEI becomes thicker and uneven, which greatly hinders the transmission of Li+ in the Cu−LTO cell. On the contrary, the rough MoS2 film can increase the effective contact area between the Cu−MoS2 collector and LTO and enhance the charge conduction, thus decreasing the internal resistance of cell and reducing the thermal effect. In addition, the MoS2 film can reduce the mismatch of the thermal expansion coefficient between the current collector and active material and effectively alleviate the thermal fatigue. Therefore, it can maintain the integrity of the LTO structure and improve the interface bond strength, thereby slowing down the formation of the SEI at the electrode surface and the interface of MoS2/LTO (Figure 13b). Therefore, the SEI in the Cu−MoS2−LTO cell is thinner and more uniform. In addition, the MoS2 film can acts as an antioxidation layer to prevent the formation of high resistance copper oxide at the interface of Cu−MoS2, promote the electron transfer between the Cu−MoS2 current collector and the active material, and prevent Li+ from diffusing to the current collector interface. Therefore, the Cu−MoS2−LTO cell exhibits excellent cycle performance and rate performance.

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.9b07269.

■

SEM image and X-ray diffraction spectrum of Li4Ti5O12; schematic diagram of semiconductor-electrochemistry model of active materials; adhesion cross-cut test of Cu−LTO, Cu−MoS2−LTO, and Cu−MoS2; cyclic voltammogram profiles of the Cu−MoS2 cell of the first three cycles at a scan rate of 0.1 mV s−1 in a potential range of 0.01−3.0 V (vs Li/Li+); SEM images of Cu−LTO electrode after 300 cycles; and SEM images of Cu−MoS2 electrode before and after the cycling (PDF)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Xin Jiang: 0000-0003-4729-8396 Notes

The authors declare no competing financial interest.

■ ■

ACKNOWLEDGMENTS This work was supported by funding from the National Natural Science Foundation of China (grant no. 51505467).

4. CONCLUSIONS In summary, we have devised an effective strategy to optimize the interface structure between the current collector and active material. This strategy is realized by magnetron sputtering technology to deposit a layer of molybdenum disulfide conductive layer on the copper foil surface, which has the advantages of simple, high efficiency, and mass production. This MoS2 conductive coating has good electronic conductivity and can increase the contact area and bond strength between the active material and the current collector, so as to improve the transport rate of electron and Li+, reduce the internal resistance of the LIBs, and effectively enhance the electrochemical performance of LIBs. Typically, after being cycled at 1 A g−1 for 300 cycles, the capacity of the Cu−LTO cell and Cu−MoS2 cell are about 114.94 and 128.35 mA h g−1, respectively, whereas the capacity of Cu−MoS2−LTO cell is as high as 373.9 mA h g−1 with a capacity retention rate of 89.1%. The MoS2 not only optimizes the interfacial structure between the current collector and LTO anode material but also provides an additional capacity contribution to the Cu− MoS2−LTO cell. Based on the SEM and XPS test analysis, we propose a dual interface model. The SEI exists on the surface of LTO and the interface between LTO and MoS2 film. Additionally, the MoS2 film can inhibit the formation of the SEI, significantly improve the charge-transfer efficiency and uniformity of the interface, prevent oxidation of the Cu foil, and thereby improve the chemical stability of the current collector. In view of the effectiveness, this strategy is equally

REFERENCES

(1) Zhong, Y.; Xia, X.; Shi, F.; Zhan, J.; Tu, J.; Fan, H. J. Transition Metal Carbides and Nitrides in Energy Storage and Conversion. Adv. Sci. 2016, 3, 1500286. (2) Mahmood, N.; Tang, T.; Hou, Y. Nanostructured Anode Materials for Lithium Ion Batteries: Progress, Challenge and Perspective. Adv. Energy Mater. 2016, 6, 1600374. (3) Liu, Z.; Chang, X.; Wang, T.; Li, W.; Ju, H.; Zheng, X.; Wu, X.; Wang, C.; Zheng, J.; Li, X. Silica-Derived Hydrophobic Colloidal Nano-Si for Lithium-Ion Batteries. ACS Nano 2017, 11, 6065−6073. (4) Kim, T.; Song, W.; Son, D.-Y.; Ono, L. K.; Qi, Y. Lithium-Ion Batteries: Outlook on Present, Future, and Hybridized Technologies. J. Mater. Chem. A 2019, 7, 2942−2964. (5) Zhang, Z.; Bao, J.; He, C.; Chen, Y.; Wei, J.; Zhou, Z. Hierarchical Carbon−Nitrogen Architectures with Both Mesopores and Macrochannels as Excellent Cathodes for Rechargeable Li−O2 Batteries. Adv. Funct. Mater. 2014, 24, 6826−6833. (6) Kang, S.-W.; Xie, H.-M.; Zhang, W.; Zhang, J.-P.; Ma, Z.; Wang, R.-S.; Wu, X.-L. Improve the Overall Performances of Lithium Ion Batteries by a Facile Method of Modifying the Surface of Cu Current Collector with Carbon. Electrochim. Acta 2015, 176, 604−609. (7) Jiang, J.; Nie, P.; Ding, B.; Wu, W.; Chang, Z.; Wu, Y.; Dou, H.; Zhang, X. Effect of Graphene Modified Cu Current Collector on the Performance of Li4Ti5O12 Anode for Lithium-ion Batteries. ACS Appl. Mater. Interfaces 2016, 8, 30926−30932. (8) Shen, L.; Shen, L.; Wang, Z.; Chen, L. In Situ Thermally Crosslinked Polyacrylonitrile as Binder for High-Performance Silicon as Lithium Ion Battery Anode. ChemSusChem 2014, 7, 1951−1956. (9) Wu, H.-C.; Lee, E.; Wu, N.-L.; Jow, T. R. Effects of current collectors on power performance of Li4Ti5O12 anode for Li-ion battery. J. Power Sources 2012, 197, 301−304. J

DOI: 10.1021/acsami.9b07269 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces (10) Reyter, D.; Rousselot, S.; Mazouzi, D.; Gauthier, M.; Moreau, P.; Lestriez, B.; Guyomard, D.; Roué, L. An Electrochemically Roughened Cu Current Collector for Si-Based Electrode in Li-Ion Batteries. J. Power Sources 2013, 239, 308−314. (11) Lee, S.; Oh, E.-S. Performance Enhancement of a Lithium Ion Battery by Incorporation of a Graphene/Polyvinylidene Fluoride Conductive Adhesive Layer Between the Current Collector and the Active Material Layer. J. Power Sources 2013, 244, 721−725. (12) Radisavljevic, B.; Radenovic, A.; Brivio, J.; Giacometti, V.; Kis, A. Single-Layer MoS2 Transistors. Nat. Nanotechnol. 2011, 6, 147− 150. (13) Wang, L.; Wang, Y.; Wong, J. I.; Palacios, T.; Kong, J.; Yang, H. Y. Functionalized MoS2 Nanosheet-Based Field-Effect Biosensor for Label-Free Sensitive Detection of Cancer Marker Proteins in Solution. Small 2014, 10, 1101−1105. (14) Portet, C.; Taberna, P. L.; Simon, P.; Flahaut, E. Modification of Al Current Collector/Active Material Interface for Power Improvement of Electrochemical Capacitor Electrodes. J. Electrochem. Soc. 2006, 153, A649−A653. (15) Wu, H.-C.; Lin, Y.-P.; Lee, E.; Lin, W.-T.; Hu, J.-K.; Chen, H.C.; Wu, N.-L. High-Performance Carbon-Based Supercapacitors Using Al Current-Collector with Conformal Carbon Coating. Mater. Chem. Phys. 2009, 117, 294−300. (16) Yang, C.-P.; Yin, Y. X.; Zhang, S. F.; Li, N. W.; Guo, Y. G. Accommodating Lithium into 3D Current Collectors with a Submicron Skeleton Towards Long-Life Lithium Metal Anodes. Nat. Commun. 2015, 6, 8058. (17) Lecoeur, C.; Tarascon, J.-M.; Guery, C. Al Current Collectors for Li-Ion Batteries Made via a Template-Free Electrodeposition Process in Ionic Liquids. J. Electrochem. Soc. 2010, 157, A641−A646. (18) Kim, H. R.; Choi, W. M. Graphene Modified Copper Current Collector for Enhanced Electrochemical Performance of Li-Ion Battery. Scr. Mater. 2018, 146, 100−104. (19) Kang, S.; Xie, H.; Zhai, W.; Ma, Z. F.; Wang, R.; Zhang, W. Enhancing Performance of a Lithium Ion Battery by Optimizing the Surface Properties of the Current Collector. Int. J. Electrochem. Sci. 2015, 10, 2324−2335. (20) Taberna, P. L.; Mitra, S.; Poizot, P.; Simon, P.; Tarascon, J.-M. High Rate Capabilities Fe3O4-Based Cu Nano-Architectured Electrodes for Lithium-Ion Battery Applications. Nat. Mater. 2006, 5, 567− 573. (21) Zhou, S.; Wang, L.; Lu, Z.; Ding, Q.; Wang, S. C.; Wood, R. J. K.; Xue, Q. Tailoring Microstructure and Phase Segregation for Low Friction Carbon-Based Nanocomposite Coatings. J. Mater. Chem. 2012, 22, 15782−15792. (22) Chen, B.; Liu, E.; He, F.; Shi, C.; He, C.; Li, J.; Zhao, N. 2D sandwich-like carbon-coated ultrathin TiO2@defect-rich MoS2 hybrid nanosheets: Synergistic-effect-promoted electrochemical performance for lithium ion batteries. Nano Energy 2016, 26, 541−549. (23) Zhang, S.; Chowdari, B. V. R.; Wen, Z.; Jin, J.; Yang, J. Constructing Highly Oriented Configuration by Few-Layer MoS2: Toward High-Performance Lithium-Ion Batteries and Hydrogen Evolution Reactions. ACS Nano 2015, 9, 12464−12472. (24) Tang, Y.; Wu, D.; Mai, Y.; Pan, H.; Cao, J.; Yang, C.; Zhang, F.; Feng, X. A Two-Dimensional Hybrid with Molybdenum Disulfide Nanocrystals Strongly Coupled on Nitrogen-Enriched Graphene via Mild Temperature Pyrolysis for High Performance Lithium Storage. Nanoscale 2014, 6, 14679−14685. (25) Dupin, J. C.; Gonbeau, D.; Martin-Litas, I.; Vinatier, P.; Levasseur, A. Amorphous Oxysulfide Thin Films MOySz (M = W, Mo, Ti) XPS Characterization: Structural and Electronic Pecularities. Appl. Surf. Sci. 2001, 173, 140−150. (26) Myamlin, V. A.; Pleskov, Y. V. The Electrochemistry of Semiconductors. Russ. Chem. Rev. 1963, 32, 207−223. (27) Zhu, X.; Zhou, S.; Yan, Q.; Wang, S. Multi-Walled Carbon Nanotubes Enhanced Superhydrophobic MWCNTs-Co/a-C:H Carbon-Based Film for Excellent Self-Cleaning and Corrosion Resistance. Diamond Relat. Mater. 2018, 86, 87−97.

(28) Ye, Y.; Liu, Z.; Liu, W.; Zhang, D.; Zhao, H.; Wang, L.; Li, X. Superhydrophobic Oligoaniline-Containing Electroactive Silica Coating as Pre-process Coating for Corrosion Protection of Carbon Steel. Chem. Eng. J. 2018, 348, 940−951. (29) Zhang, B.; Liu, Y.; Huang, Z.; Oh, S.; Yu, Y.; Mai, Y.-W.; Kim, J.-K. Urchin-Like Li4Ti5O12-Carbon Nanofiber Composites for High Rate Performance Anodes in Li-Ion Batteries. J. Mater. Chem. 2012, 22, 12133−12140. (30) Yang, L.; Wang, S.; Mao, J.; Deng, J.; Gao, Q.; Tang, Y.; Schmidt, O. G. Hierarchical MoS2/Polyaniline Nanowires with Excellent Electrochemical Performance for Lithium-Ion Batteries. Adv. Mater. 2013, 25, 1180−1184. (31) Fang, X.; Yu, X.; Liao, S.; Shi, Y.; Hu, Y.-S.; Wang, Z.; Stucky, G. D.; Chen, L. Lithium Storage Performance in Ordered Mesoporous MoS2 Electrode Material. Microporous Mesoporous Mater. 2012, 151, 418−423. (32) Haetge, J.; Hartmann, P.; Brezesinski, K.; Janek, J.; Brezesinski, T. Ordered Large-Pore Mesoporous Li4Ti5O12 Spinel Thin Film Electrodes with Nanocrystalline Framework for High Rate Rechargeable Lithium Batteries: Relationships among Charge Storage, Electrical Conductivity, and Nanoscale Structure. Chem. Mater. 2011, 23, 4384−4393. (33) Fang, X.; Guo, X.; Mao, Y.; Hua, C.; Shen, L.; Hu, Y.; Wang, Z.; Wu, F.; Chen, L. Mechanism of Lithium Storage in MoS2 and the Feasibility of Using Li2S/Mo Nanocomposites as Cathode Materials for Lithium-Sulfur Batteries. Chem.Asian J. 2012, 7, 1013−1017. (34) Cai, L.; Zhang, Q.; Mwizerwa, J. P.; Wan, H.; Yang, X.; Xu, X.; Yao, X. Highly Crystalline Layered VS2 Nanosheets for All-Solid-State Lithium Batteries with Enhanced Electrochemical Performances. ACS Appl. Mater. Interfaces 2018, 10, 10053−10063. (35) Cook, J. B.; Kim, H.-S.; Yan, Y.; Ko, J. S.; Robbennolt, S.; Dunn, B.; Tolbert, S. H. Mesoporous MoS2 as a Transition Metal Dichalcogenide Exhibiting Pseudocapacitive Li and Na-Ion Charge Storage. Adv. Energy Mater. 2016, 6, 1501937. (36) Chao, D.; Zhu, C.; Yang, P.; Xia, X.; Liu, J.; Wang, J.; Fan, X.; Savilov, S. V.; Lin, J.; Fan, H. J. Array of Nanosheets Render Ultrafast and High-Capacity Na-Ion Storage by Tunable Pseudocapacitance. Nat. Commun. 2016, 7, 12122. (37) Kalimuldina, G.; Taniguchi, I. Electrochemical Properties of Stoichiometric CuS Coated on Carbon Fiber Paper and Cu Foil Current Collectors as Cathode Material for Lithium Batteries. J. Mater. Chem. A 2017, 5, 6937−6946. (38) Kalimuldina, G.; Taniguchi, I. Synthesis and electrochemical characterization of stoichiometric Cu2S as cathode material with high rate capability for rechargeable lithium batteries. J. Power Sources 2016, 331, 258−266. (39) Kalimuldina, G.; Taniguchi, I. High Performance Stoichiometric Cu2S Cathode on Carbon Fiber Current Collector for Lithium Batteries. Electrochim. Acta 2017, 224, 329−336. (40) Zhuang, W.; Fan, X.; Li, W.; Li, H.; Zhang, L.; Peng, J.; Cai, Z.; Mo, J.; Zhang, G.; Zhu, M. Comparing Space Adaptability of Diamond-Like Carbon and Molybdenum Disulfide Films Toward Synergistic Lubrication. Carbon 2018, 134, 163−173. (41) Li, H.; Zhang, G.; Wang, L. Low Humidity-Sensitivity of MoS2/ Pb Nanocomposite Coatings. Wear 2016, 350−351, 1−9.

K

DOI: 10.1021/acsami.9b07269 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX